Phase transformation induced high-strength titanium carbide reinforced stainless steel composites with stable thermo-mechanical properties for high temperature applications

doi:10.21203/rs.3.rs-2000232/v1

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Phase transformation induced high-strength titanium carbide reinforced stainless steel composites with stable thermo-mechanical properties for high temperature applications

https://doi.org/10.21203/rs.3.rs-2000232/v1

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The high-temperature properties such as the tensile strength, hardness, coefficient of thermal expansion (CTE), and thermal conductivity of titanium carbide (TiC)-reinforced stainless steel (SS) matrix composites fabricated by the infiltration process were investigated at various temperatures. Highly concentrated, closely adjacent TiC particulate reinforcement effectively suppressed dramatic variations in the CTE and thermal conductivity properties originating from the phase transformation of the SS431 matrix with superior mechanical properties. The strengthening effect of the addition of TiC increased as the temperature increased and showed a higher value after the phase transformation. The TiC–SS431 composite showed great applicability for advanced high-temperature-resistant structural materials.

TiC

particulate reinforcement

SS431

phase transformation

high-temperature steel

structural materials

Superior mechanical, thermal, and thermomechanical properties are of fundamental interest for high-temperature applications that require materials to retain their properties and shapes at different temperatures. However, with increasing temperature, the physical and mechanical properties of materials change, resulting in degradation of material properties and shortened reliability of structures ^{[1, 2]}. Ferritic steels contain several phases, including ferrite, pearlite, and martensite, all of which have body-centered cubic (BCC) structures. In other words, there are atoms at the eight corners of a cube and one at the center of the cube. Pure iron atoms were arranged in a body-centered cubic structure up to 911°C. Above this temperature, the structure undergoes a phase transition to face-centered cubic (FCC). The transition from BCC to FCC structure results in an 8 to 9% increase in density of iron and steel, resulting in volumetric shrinkage in size as material is heated above the transition temperature ^[3–5]. When phase transformation occurs at high temperatures, the dimensional stability of iron and steels deteriorates significantly. Hence, efforts have been directed toward improving the high-temperature characteristics of steels ^[6–8].

Owing to their superior physical and mechanical properties, metal matrix composites (MMCs) can effectively enhance the performance limits of conventional metals ^[9–16]. Although MMCs are promising materials for the chemical and process industries, little attention has been paid to their properties owing to their high melting temperature, poor wettability, and density mismatch between the reinforcement and metal matrix ^{[10, 17]}. Typically, steel matrix composites are reinforced with ceramic particles, such as carbides (titanium carbide (TiC), SiC, WC), borides (TiB₂), and oxides (Al₂O₃, ZrO₂) ^[18–29]. TiC particles can be considered as suitable reinforcements for steel matrix composites owing to their high hardness, low density, good wettability with steel, chemical inertness, and high melting point ^[23–30]. To date, most investigations of TiC-reinforced steel composites have focused on room-temperature mechanical properties such as hardness ^[23], compressive strength ^[24], transverse rupture strength ^[28], and wear properties ^[31]. The special properties of this material have expanded its application to high-temperature components, for which superior strength and stable thermal and thermo-mechanical properties are essential. However, comparatively little research has been conducted on their high-temperature properties. Furthermore, the effect of solid-state phase transformation of the steel matrix on the high-temperature properties of TiC-reinforced steel composites has not been reported.

In this study, the effects of TiC reinforcement on the high-temperature mechanical, thermomechanical, and thermal properties of steel matrix composites were evaluated using SS431, as well as using a TiC–SS431 composite fabricated by the gas pressure infiltration (GPI) process while considering the phase transformation of the steel matrix.

TiC powders (Changsha Langfeng Metal Materials Co., Ltd, 2500 mesh, China) and stainless steel (SS)431 (SeAH css Corp., Republic of Korea), which is a high-chromium (Cr) low-nickel (Ni) stainless steel with high strength and good corrosion resistance, were used as the raw materials. The nominal chemical composition of SS431 used in this study was Fe-(15–17)Cr-(1.25 ~ 2.5)Ni-1Si-1Mn-0.2C. To fabricate the TiC–SS431 composites, TiC powders were uniaxially pressed in a steel die at 6 MPa. After compaction, pressed bulk TiC was sintered at 1200°C for 60 min under an Ar atmosphere to fabricate weakly bonded TiC preforms with a very porous structure. The open porosity of the fabricated cylindrical TiC preform (diameter: 50 mm; height: 25 mm), calculated by measuring its weight and size, was approximately 40%. TiC-reinforced SS431 composites were fabricated by the GPI process, which is capable of temperatures up to 2000°C and can apply a maximum pressure of 30 bar. The GPI process was developed to infiltrate molten metal into a ceramic preform via hydrostatic pressure, leading to a uniform dispersion of the reinforcement inside the steel matrix. SS431 was melted under a vacuum of 10^− 2 Torr in a chamber. The TiC preform was separately heated to 1600°C in a MgO mold in the same chamber. The molten SS431 alloy was infiltrated into the preheated TiC preform and then pressurized with an Ar pressure of 10 bar. No further thermal treatments were performed in this study.

The particle size of the as-received TiC powder was analyzed using a laser particle size analyzer (LS 13 320, Beckman Coulter). The microstructures of the composites were investigated in detail by scanning electron microscopy (SEM; JSM-6610LV, JEOL) and transmission electron microscopy (TEM; JEM 2100F, JEOL). The TEM samples were prepared using a focused ion beam (FIB) system. The thermal conductivity of the composite was calculated using the expression λ = α ρ C_p, where α is the thermal diffusivity, ρ is the bulk density, and C_p is the specific heat. The bulk densities of the composites were measured using a method based on Archimedes’ law. The thermal diffusivity and specific heat of the SS431 and TiC–SS431 composites were measured at about 25°C using a laser flash apparatus (LFA 457, Netzsch). To obtain reliable results, five tests were performed for each sample, and the mean value obtained. Coefficient of thermal expansion (CTE) values of SS431 and of the TiC–SS431 composite were also measured with heating rate of 5°C/min from RT to 900°C using a thermomechanical analyzer (TMA 402 F3, Netzsch). The phase of the TiC–SS431 composite at various temperatures (RT–900°C) was evaluated using a high-temperature X-ray diffractometer (Ultima IV, Rigaku) with Cu Kα radiation at 40 kV/40 mA. The Vickers hardness values from RT to 800°C for the SS431 and TiC–SS431 composites were measured using a high-temperature Vickers hardness testing system (AVK-HF, Mitutoyo) in an Ar environment. The measured load was 5 Kg_f and the holding time was 11–12 s. Each specimen was measured three times. High-temperature tensile tests were carried out using a universal testing machine (5882 model, INSTRON) equipped with a 100 kN load cell. Tensile tests were conducted at 700, 800, and 900°C at a strain rate of 0.005/s. The specimens were machined by wire electrode cutting with a cross-section of 6.5 × 1.5 mm² and a gauge length of 5 mm.

3.1. Microstructural characterization of TiC–SS431composite

Figure 1 shows the microstructures of the TiC powder, TiC preform, and TiC–SS431 matrix composites. Figure 1a shows an SEM image of the TiC powder used in this study as a reinforcement. The as-received TiC powder exhibited an irregular faceted shape. The results showed a bimodal distribution, with small particles of approximately 1 µm and a greater amount of larger particles of approximately 10 µm. Measured particle size distribution D50 is 13.67 µm as shown in the inset of the Fig. 1a. Figure 1b shows the various shapes of TiC preforms. TiC powders and preforms showed similar microstructures fabricated by low-temperature sintering at 1200°C. However, the amount of small particles were slightly reduced compared to the as-received TiC powder. The preform was maintained in the form of powder without severe grain growth through low-temperature sintering. The TiC particles of the preform were weakly attached to each other while maintaining the preform shape. Figure 1c shows an SEM image of the TiC–SS431 composite fabricated by the GPI process at 1600°C. The TiC particles were uniformly dispersed in the SS431 matrix. The measured average volume fraction of TiC obtained using ten low-magnification SEM images was approximately 58%. The observed TiC reinforcements in the SS431 matrix had relatively round shapes compared with the as-received TiC particles because of the partial dissolution and reprecipitation of TiC ^{[28, 29]}. In addition, the TiC–SS431 sample images showed that the TiC preform did not break and maintained its shape during the infiltration process.

To gain further insight into the microstructure of the TiC–SS431 composites, TEM images and the corresponding EDS maps were obtained, as shown in Fig. 2. In the low-magnification TEM image (Fig. 2a), some TiC particles appear to be attached to each other; however, in the high-magnification TEM image, molten Fe can be seen to infiltrate into narrow regions between closely adjacent TiC particles (Fig. 2c). Precipitates, such as Cr carbide, were not observed at the TiC/SS431 interface, as shown in TEM-EDS mapping images (Fig. 2b). The SS431 melt was well infiltrated into a very narrow region of approximately 170 nm between TiC particle and TiC particle, without interfacial defects, owing to the good wettability between TiC and the SS431 matrix with the aid of the applied gas pressure (Fig. 2c). In the high-resolution TEM observation, a semi-coherent interface was found along the low-index plane, such as the {111} plane of TiC and the {110} plane of SS431 (Fig. 2d). Selected area diffraction pattern (SADP) analysis confirmed that the {111} plane of TiC was distorted by approximately 1° toward the {110} plane of SS431. From this interfacial relationship, it was found that the interfacial bonding strength between TiC and the SS431 matrix was excellent.

3.2. Phase characterization of TiC–SS431composite

In this study, the phase diagram of SS431 under TiC saturation was calculated based on the solubility of TiC in Fe, as shown in Fig. 3. Thermodynamic calculations were performed using FactSage 7.3 thermo-chemical software and FSstel database. The calculated saturation content of TiC at 1600°C, which is the fabrication temperature of the TiC–SS431 composite, was 6.8595 wt.%. The saturation content of TiC decreased by 4.2824 wt.% at 1500°C. The phase diagram of SS431 under TiC saturation conditions revealed that the phase transformation started at 657.53°C (Ac₁) and finished at 830.6°C (Ac₃), as shown in Fig. 3. Notably, the phase transformation of the TiC–SS431 composite was mainly generated from about 750°C. The calculation of SS431 with and without TiC revealed that the effect of TiC on the transformation temperature was negligible, with a difference of less than 1°C.

The phases of the TiC–SS431 composite at various temperatures were analyzed using high–temperature XRD to determine the transformation temperature of the SS431 matrix and the microstructural changes in the TiC–SS431 composite. Figure 4 shows the high-temperature XRD results of the TiC–SS431 composite in the temperature range RT–900°C. The XRD pattern of the TiC–SS431 composite measured at RT showed the presence of α-Fe (ferrite), γ-Fe (austenite), and TiC phases without impurity peaks. The presence of weak γ-Fe (FCC) peaks suggests that the partial dissolution of TiC particles may have resulted in an increase in the carbon concentration of the TiC–SS431 composites, leading to an increase in residual austenite. The XRD pattern of the TiC–SS431 composites measured at 900°C clearly revealed the formation of a γ-Fe (FCC) phase with a decreased peak intensity of the α-Fe phase. The XRD results agreed well with the FactSage calculation results. It can be seen that TiC and Fe peaks move to a lower angle as the temperature increases. This is attributed to the expansion of the interatomic distance caused by increased atomic thermal vibrations as the temperature increases. However, this phenomenon started to slow from 800°C, as can be seen in the TiC peaks, and the XRD pattern measured at 900°C shows almost the same 2θ position of TiC as that measured at 800°C, indicating suppression of the expansion of TiC interatomic distance by shrinkage of the Fe matrix due to phase transformation.

3.3. Thermal expansion mechanism of TiC–SS431composite

To understand the effects of TiC on the thermal expansion characteristics near the phase transformation temperature of steel, CTE measurements were performed. Figure 5 shows the thermal expansion properties of the SS431 and TiC–SS431 composites. Figure 5a shows the length changes (dL/L₀×10³) of the SS431 and TiC–SS431 composites as a function of temperature (L₀ is the initial length and dL is the elongation). Both SS431 and TiC–SS431 composites show linear length increases with increasing temperature up to 700°C. However, the increased length of SS431 suddenly decreased over 775–830°C and then increased again with increasing temperature. This phenomenon has been commonly reported for various steels for the BCC to FCC phase transition with an increased packing density of Fe atoms, and agrees well with thermodynamic calculations. Meanwhile, there was no dramatic length reduction of the TiC–SS431 composite at temperatures above 700°C. An almost linear increase in length can be observed in the graph up to 900°C without the effect of matrix length reduction. This trend was also observed during the cooling process. Figure 6c shows the measured CTE values for the SS431 and 58 vol.% TiC–SS431 composites at temperatures up to 700°C. Measured CTE values of the TiC–SS431 composite were 21.9–23.6% lower than those of SS431 over the entire temperature range. The measured CTE values of the TiC–SS431 composites from 25 to 100°C were compared with the calculated CTE values as a function of the TiC volume fraction using the rule of mixtures (ROM), Turner’s model ^[32], and Kerner’s model ^[33], as shown in Fig. 3d. The expressions used to calculate ROM, Tuner’s model, and Kerner’s model are presented in Table 1, where α is the coefficient of thermal expansion, V is the volume fraction, K is the bulk modulus, G is the shear modulus, and E is the Young’s modulus. The subscripts c, f, and m represent the composite, filler, and matrix, respectively.

Table 1

Model for coefficient of thermal expansion of composites
Models	Expressions
ROM	${\alpha }_{c}=\left(1-{V}_{f}\right){\alpha }_{m}+{V}_{f}{\alpha }_{f}$
Turner's Model	${\alpha }_{c}=\frac{\left(1-{V}_{f}\right){\alpha }_{m}{K}_{m}+{V}_{f}{\alpha }_{f}{K}_{f}}{\left(1-{V}_{f}\right){K}_{m}+{V}_{f}{K}_{f}}$
Kerner's Model	${\alpha }_{c}=\left(1-{V}_{f}\right){\alpha }_{m}+{V}_{f}{\alpha }_{f}+\frac{{V}_{f}\left(1-{V}_{f}\right)\left({\alpha }_{f}-{\alpha }_{m}\right)\left({K}_{f}-{K}_{m}\right)}{\left(1-{V}_{f}\right){K}_{m}+{V}_{f}{K}_{f}+3{K}_{f}{K}_{m}/4{G}_{m}}$

The results revealed that the measured CTE value (9.42 ppm/°C) of the 58 vol.% TiC–SS431 composite was very similar to that of the Kerner’s model. Turner’s and Kerner’s models can be used to theoretically calculate the coefficients of the CTE of composites. Kerner’s model considers the shear effects at the boundaries between particles and the matrix, whereas Turner’s model does not account of such effects. Consequently, Kerner’s model has been widely accepted for the theoretical evaluation of the CTEs of composites ^[38]. The good agreement between the experimental and calculated values indicate good interfacial bonding between TiC and the SS431 matrix. Figure 5d shows the length changes of the SS431 and TiC–SS431 composites (solid lines) and the length values calculated by ROM with various TiC volume fractions (dashed lines) in the temperature range of 500–900°C. The calculated length changes revealed that the length shrinkage of the TiC–SS431 composite decreased with increasing TiC content. In the case of 60 vol.% TiC–SS431 composite, a reduction in length was clearly observed at about 780°C, whereas almost no reduction in length was observed for the 58 vol.% TiC–SS431 composite. In the case of the ROM model and the uniformly dispersed composites (Fig. 5e), the length change due to the phase transformation occurred at 60 vol.% TiC–SS431 composite. Because the highly concentrated, closely adjacent TiC particles with a semi-coherent interface with the SS431 matrix maintained the entire structure of the TiC–SS431 composite fabricated by the infiltration process, as shown in Fig. 2, shrinkage of the composite was suppressed during phase transformation. Figure 5f shows a schematic of the TiC–SS431 composite, which sustains its structure during the phase transformation (α-Fe → γ-Fe) of the Fe matrix by highly concentrated, closely adjacent TiC reinforcements at high temperatures.

3.4. Thermal conductivity of TiC–SS431composite

Figure 6 and Table 2 show the thermal conductivities of the SS431 and TiC–SS431 composites measured at various temperatures. SS431 exhibited a relatively higher thermal conductivity than that of the TiC–SS431 composite up to 500°C (Fig. 6a). However, the thermal conductivity of SS431 dramatically decreased at approximately 700°C, largely due to the decreased thermal diffusivity; the thermal conductivity then remained below the value of the TiC–SS431 composite. The decreased thermal conductivity of SS431 increased again with an increase in the temperature to 900°C. Because the FCC structure commonly has a higher conductivity than the BCC structure owing to its higher packing density of metal atoms, the thermal diffusivity and thermal conductivity of SS431 at 900°C increased again with phase transformation (BCC → FCC). The thermal conductivity of the TiC–SS431 composite increased almost linearly with increasing temperature, regardless of the phase transformation of the Fe matrix. Figure 6b shows the measured thermal diffusivities and specific heats of the SS431 and TiC–SS431 composite. The specific heat of SS431 increased linearly with increasing temperature, whereas the thermal diffusivities of SS431 decreased with increasing temperature up to 700°C because of the decreased mean free path of free Fe electrons by lattice vibration at high temperature. This phenomenon occurred severely at high temperatures, resulting in the decreased thermal conductivity of SS431 at 700°C. Consequently, the thermal conductivity of the TiC–SS431 composite gradually increased to 500°C owing to the increased specific heat and decreased at higher temperatures because of the significantly decreased thermal diffusivity. However, as the phase transformation occurred, the thermal diffusivity increased at approximately 900°C, resulting in an increased thermal conductivity. Therefore, the results reveal that the phase transformation of the Fe matrix does not affect the specific heat but only the thermal diffusivity. In contrast, the measured specific heat of the TiC–SS431 composite increased linearly with increasing temperature. Unlike SS431, the thermal diffusivities of the TiC–SS431 composite did not severely decrease with increasing temperature and maintained values of over 5 mm²s^− 1 because of the highly concentrated TiC reinforcement with a semi-coherent interface with the SS431 matrix. The calculated interfacial thermal conductance of TiC/SS431 based on the Hasselman–Johnson model ^[35] was approximately 1–2×10⁷ W m^− 2K^− 1. The Hasselman–Johnson model is expressed as:

$${\lambda }_{eff}={\lambda }_{m}\frac{2\left(\frac{{\lambda }_{d}}{{\lambda }_{m}}-\frac{{\lambda }_{d}}{a{h}_{c}}-1\right){V}_{d}+\frac{{\lambda }_{d}}{{\lambda }_{m}}+\frac{2{\lambda }_{d}}{a{h}_{c}}+2}{\left(1-\frac{{\lambda }_{d}}{{\lambda }_{m}}+\frac{{\lambda }_{d}}{a{h}_{c}}\right){V}_{d}+\frac{{\lambda }_{d}}{{\lambda }_{m}}+\frac{2{\lambda }_{d}}{a{h}_{c}}+2}$$

where λ_eff is the effective thermal conductivity, λ_m and λ_d are respectively the thermal conductivities of the matrix and inclusion, V_d is the volume fraction of the inclusion, a is the radius of the inclusion, and h_c is the interface thermal resistance.

Table 2

Thermal conductivities of SS431 and TiC–SS431 composites measured at various temperatures.
		Testing temperature (°C)
		25	300	500	700	900
SS431 (7.691 g/cm³)	Thermal diffusivity (cm²s^− 1)	0.05941	0.05537	0.04817	0.03288	0.05685
	Specific heat (J g^− 1K^− 1)	0.456	0.573	0.730	0.824	0.884
	Thermal conductivity (Wm^− 1K^− 1)	20.835	24.4	27.032	20.837	38.649
TiC-SS431 (5.812 g/cm³)	Thermal diffusivity (cm²s^− 1)	0.05556	0.05052	0.05131	0.05142	0.05751
	Specific heat (J g^− 1K^− 1)	0.527	0.679	0.739	0.781	0.818
	Thermal conductivity (Wm^− 1K^− 1)	17.013	19.936	22.035	23.339	27.343

3.5. Mechanical properties of TiC–SS431composite at elevated temperatures

The high-temperature mechanical properties of SS431 and TiC–SS431 composite were evaluated. Figure 7a shows the Vickers hardness of the SS431 and TiC–SS431 composites as a function of temperature. The Vickers hardness of SS431 gradually decreased with increasing temperature. The Vickers hardness of the TiC–SS431 composite was significantly higher than that of SS431 over the entire temperature range. In particular, the hardness of the TiC–SS431 composite was 3.65 times higher than that of SS431 at 800°C. Gradual shrinkage by the phase transformation of SS431 during isothermal measurement at 800°C with the aid of applied pressure led to a decrease in the Vickers hardness of SS431. The effects of pressure on the phase transformations of various materials have been reported in previous studies ^{[36, 37]}. Because the highly concentrated, closely adjacent TiC particles sustained the entire structure of the TiC–SS431 composite during the phase transformation of the matrix, the strengthening effect of TiC increased at 800°C as compared with the values measured below 800°C. This result is in good agreement with the CTE results for the TiC–SS431 composite.

To evaluate the high-temperature strengths of SS431 and TiC–SS431 composites, tensile tests were performed at 700, 800, and 900°C. The tensile stress-strain curves of the SS431 and TiC–SS431 composites are shown in Figs. 8a–8c. The high-temperature tensile strengths of the TiC–SS431 composites were significantly improved compared with those of the SS431 alloy, regardless of the temperature. The average strengthening ratios (σ_c/ σ_m) at 700°C (matrix: BCC), 800°C (matrix: BCC + FCC), and 900°C (matrix: FCC) were respectively 2.77, 3.85, and 4.61, as shown in Fig. 8d and Table 2. Where, σ_c is strength of the composite, and σ_m is strength of the matrix. These results indicate that the strengthening effect of TiC increases as the testing temperature increases. While the strength of SS431 decreased with increasing test temperature, uniformly dispersed TiC, which has a high melting point and hardness, sustained the structural strength of the composite at the tested temperatures. Figure 8e shows the fracture morphologies of the TiC–SS431 composite after tensile testing at various temperatures. The inset of Fig. 8e shows the longitudinal cross sections after tensile testing. SEM images of the TiC–SS431 composites tested at 700 and 800°C show a majority of TiC cleavage fractures with no interfacial debonding between TiC and the steel matrix, indicating excellent bonding between TiC and the SUS431 matrix. These results indicate that initial failure occurred in TiC when the load transferred from the matrix to TiC exceeded the critical strength of TiC. The TiC–SS431 composites tested at 900°C were fractured in a mixed mode with ductile rupture of the SUS431 matrix and cleavage fracture of TiC reinforcements. Fractured TiC particles were rarely observed near the fractured surface of the TiC–SS431 composites tested at 700 and 800°C. On the other hand, multiple cracks with fractured TiC particles, indicating suppression of local stress concentration, were observed in an area 350 µm away from the fractured surface of TiC–SS431 composites tested at 900°C. This is attributed to the increased resistance of the TiC–SS431 composites to fracture at 900°C. The present investigation involves an additional hardening effect induced by the phase transformation of the SS431 matrix. Because the coherency of the TiC/SS431 interface is expected to be maintained after the phase transformation, a large amount of stress should be generated at the interface. It can thus be inferred from the results of the suppressed shrinkage of the matrix that the strain energy at and near the TiC interface increases significantly. As the interfacial strain energy increased, the interfacial strengthening effect increased with the effective dispersion of the applied stress in the composite. Therefore, the strengthening effect of TiC increases as the temperature increases above the FCC transformation temperature, as shown in Fig. 8d and Table 3.

Table 3

Tensile strengths of SS431 and TiC–SS431 composites measured at high temperature.
Specimen	Tensile Strength (MPa)
Specimen	700°C	800°C	900°C
SS431	186	127	94
TiC–SS431 composite	512 519	486 491	401 465
K (σ_c/ σ_m)	2.75 2.79	3.82 3.86	4.26 4.94

Highly concentrated, closely adjacent TiC–SS431 composites were successfully fabricated using the GPI process. While SS431 showed large changes in thermal conductivity and CTE at high temperatures because of the phase transformation from BCC to FCC, the TiC–SS431 composite had stable thermal conductivity and CTE values regardless of phase transformation. Furthermore, the addition of TiC significantly improves the tensile strength and hardness of SS431 at high temperatures. The strengthening effect increased as the temperature increased owing to the applied compressive stress at TiC after the phase transformation of SS431, leading to effective load transfer from the SS431 matrix to TiC. These results indicate the possibility of exploiting TiC as a contributing reinforcement for improving the mechanical and physical properties of steel matrix composites at high temperatures. Therefore, the highly concentrated TiC–SS431 composite is expected to be applicable as an advanced high-temperature-resistant structural material.

Acknowledgment

This work was financially supported by the National Research Foundation of Korea (NRF) grant funded by the Korea government (MSIP) (2020M3H4A3105943, 2022M3H4A3046292).

Author Contributions Statement

S. Cho: writing—original draft, funding acquisition, project administration. J. Kim: data curation, software. J. Lee: methodology, formal analysis. S. Shin: formal analysis, visualization. D. Lee: investigation, formal analysis. M. Kim: data curation. H.-U. Hong: validation, methodology. M. Choi: visualization. Y.-S. Lee: conceptualization, writing—review and editing. I. Jo: conceptualization, writing—review and editing. S.-K. Lee: supervision.

Declaration of Competing Interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

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No competing interests reported.

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Phase transformation induced high-strength titanium carbide reinforced stainless steel composites with stable thermo-mechanical properties for high temperature applications

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Abstract

Figures

1. Introduction

2. Experimental Methods

3. Results And Discussion

3.1. Microstructural characterization of TiC–SS431composite

3.2. Phase characterization of TiC–SS431composite

3.3. Thermal expansion mechanism of TiC–SS431composite

3.4. Thermal conductivity of TiC–SS431composite

3.5. Mechanical properties of TiC–SS431composite at elevated temperatures

4. Conclusion

Declarations

References

Additional Declarations

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Models	Expressions
ROM	\({\alpha }_{c}=\left(1-{V}_{f}\right){\alpha }_{m}+{V}_{f}{\alpha }_{f}\)
Turner's Model	\({\alpha }_{c}=\frac{\left(1-{V}_{f}\right){\alpha }_{m}{K}_{m}+{V}_{f}{\alpha }_{f}{K}_{f}}{\left(1-{V}_{f}\right){K}_{m}+{V}_{f}{K}_{f}}\)
Kerner's Model	\({\alpha }_{c}=\left(1-{V}_{f}\right){\alpha }_{m}+{V}_{f}{\alpha }_{f}+\frac{{V}_{f}\left(1-{V}_{f}\right)\left({\alpha }_{f}-{\alpha }_{m}\right)\left({K}_{f}-{K}_{m}\right)}{\left(1-{V}_{f}\right){K}_{m}+{V}_{f}{K}_{f}+3{K}_{f}{K}_{m}/4{G}_{m}}\)